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Fabrication and thermoelectricproperties of Yb-doped ZrNiSn half-Heusler alloysXiao-Hua Liu a , Jian He b , Han-Hui Xie a , Xin-Bing Zhao a & Tie-Jun Zhu aa State Key Laboratory of Silicon Materials and Department ofMaterials Science and Engineering, Zhejiang University, Hangzhou,310027, P.R. Chinab Department of Physics and Astronomy, Clemson University, SouthCarolina, 29634-0978, USAVersion of record first published: 16 Dec 2011.

To cite this article: Xiao-Hua Liu , Jian He , Han-Hui Xie , Xin-Bing Zhao & Tie-Jun Zhu (2012):Fabrication and thermoelectric properties of Yb-doped ZrNiSn half-Heusler alloys, InternationalJournal of Smart and Nano Materials, 3:1, 64-71

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International Journal of Smart and Nano MaterialsVol. 3, No. 1, March 2012, 64–71

Fabrication and thermoelectric properties of Yb-doped ZrNiSnhalf-Heusler alloys

Xiao-Hua Liua, Jian Heb, Han-Hui Xiea, Xin-Bing Zhaoa and Tie-Jun Zhua*

aState Key Laboratory of Silicon Materials and Department of Materials Science and Engineering,Zhejiang University, Hangzhou 310027, P.R. China; bDepartment of Physics and Astronomy,

Clemson University, South Carolina 29634-0978, USA

(Received 14 August 2011; final version received 1 November 2011)

Half-Heusler (HH) alloys constitute an important class of materials that exhibitpromising potential in high-temperature thermoelectric (TE) power generation. In thiswork, we synthesized Zr1−xYbxNiSn (x = 0, 0.01, 0.02, 0.04, 0.06 and 0.10) HH alloysusing a time-efficient levitation melting and spark plasma sintering procedure. X-raydiffraction showed that the samples were predominantly single phased, and that thelattice constant increased systematically with increasing Yb doping ratio. The dopingeffects of Yb on the thermoelectric properties were studied. It was found that Yb dop-ing consistently decreased the electrical and thermal conductivities. On the other hand,the effects of Yb doping on the Seebeck coefficient were found to be non-monotonic.The magnitude of the Seebeck coefficient (n-type) was increased upon Yb doping up tox = 0.02, above which Yb doping introduced notable p-type conduction. As a result, theroom-temperature Seebeck coefficient of the x = 0.10 sample became positive althoughthe magnitude was not high. The thermoelectric figure of merit, ZT , reached a maxi-mum of ∼0.38 at 900 K for the x = 0.01 sample. Selective doping on the Ni and Snsites are necessary to further optimize the TE performance of Zr1−xYbxNiSn alloys.

Keywords: half-Heusler alloy; thermoelectric power; levitation melting; spark plasmasintering; Yb doping

1. Introduction

In recent years the global energy crisis has resulted in a pressing need for the develop-ment of high-performance thermoelectric (TE) materials for direct thermal-to-electricalpower generation. The performance of a TE material is gauged by the dimensionless figureof merit, ZT = α2σT/κ , where α is the Seebeck coefficient, σ the electrical conduc-tivity, κ the thermal conductivity and T the absolute temperature. In the simplest case,κ = κL + κe, where κL and κe are the lattice and carrier thermal conductivities, respec-tively. The power factor, α2σ , which is intimately related to the electronic band structure,is in general optimized via doping to maximize the ZT [1,2].

Half-Heusler (HH) ternary intermetallic compounds are formulated as XYZ, where Xand Y are two different transition metal elements, and Z is a sp element such as Sn, Sbor Bi. HH compounds are of considerable interest for TE power generation above roomtemperature. HH alloys adopt an MgAgAs-type structure (space group F43m), consisting

*Corresponding author. Email: zhutj@zju.edu.cn

ISSN 1947-542X online© 2012 Taylor & Francishttp://dx.doi.org/10.1080/19475411.2011.637994http://www.tandfonline.com

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of three interpenetrating face centered cubic (fcc) sublattices. The crystallographic sites(0, 0, 0) and (1/4, 1/4, 1/4) are occupied by two different transition metal elements X andY, respectively, the (1/2, 1/2, 1/2) site is occupied by a sp element Z, whereas the site(3/4, 3/4, 3/4) is vacant [3]. The band structure of HH alloys is profoundly affected by thevalence electron count (VEC): the compounds are narrow band gap semiconductors whenVEC = 8 or 18, while the deviation from VEC = 18 may lead to metallic conduction [4].

From the thermoelectrics point of view, the group of HH alloys most intensively inves-tigated is XNiSn (X = Ti, Zr, Hf). The XNiSn alloys are narrow band gap (0.1–0.2 eV)semiconductors [5,6]. The narrow band gaps give rise to a large effective mass that in turnleads to a large Seebeck coefficient [7]. A large Seebeck coefficient ∼100 µV K−1 and ahigh electrical conductivity ∼103–104 Scm−1 have been attained [8,9]. In addition, XNiSnalloys have high melting points of 1400–1600 K [10] and excellent chemical stability withnegligible sublimation near their melting points.

The TE performance of these XNiSn alloys can be optimized for practical TE powergeneration via selectively doping the three filled sublattices [11]. For example, dopingon the Sn site tunes carrier concentration [12–14], while substituting the Ti and Ni sitespromotes the mass fluctuations and strain fields that help reduce the lattice thermal con-ductivity [15–18]. The ZT of n-type Hf0.6Zr0.4NiSn0.98Sb0.02 is ∼1 at 1000 K, which isthe best and reproducible TE performance that n-type HH TE materials have attained sofar [12].

As far as making a practical TE power generation module is concerned, both n- andp-type materials are needed, better with similar composition and mechanical properties.In contrast to the promising TE properties of n-type HH alloys, p-type HH compoundsare rarely reported, and the reported TE performance was not satisfactory [19–22]. Forexample, ZT values of ∼0.5 and ∼0.29 were attained in the p-type Zr0.5Hf0.5CoSb0.8Sn0.2

[23] and ErNiSn0.01Sb0.99 [19], respectively. It is thus highly desirable to control the typeand the number of charge carriers in XNiSn by doping on the X, Ni or Sn sites [7,24,25].

In this vein, Katsuyama et al. [24] and Hlil et al. [26] proposed that Y doping inZrNiSn was equivalent to doping with acceptors since the number of valence electronsof Y (4d15s2) is one less than that of Zr (4d25s2), so heavy Y doping might increase thehole concentration and even turn ZrNiSn into p-type. Zhu et al. [27] recently showed that Ydoping in Hf0.6Zr0.4NiSn0.98Sb0.02 reduced thermal conductivity due to the reduced carrierthermal conductivity, and increased the Seebeck coefficient due to the decreased carrierconcentration.

Another promising class of dopants is the rare earth lanthanide elements [28].Lanthanide elements are characterized by 4f electrons. It is established that rare earth-containing compounds often have strong hybridization between the 4f electrons and theconduction band of the parent compounds, giving rise to enhanced electronic densityof states near the Fermi energy and thus higher Seebeck coefficient. Moreover, heat ismostly carried by short-wavelength phonons at temperatures where HH TE materialsoperate, and hence the rare earth doping will introduce short-range disorder and strainfields to effectively scatter heat-carrying phonons. Nonetheless, the electrical conductiv-ity, Seebeck coefficient and thermal conductivity are inter-related in a given material, thedetailed effects of rare earth doping on the TE properties are material specific and subjectto exploratory experimental investigations.

In order to verify whether and to what extent the rare earth lanthanide doping canchange the carrier concentration and even the type of the primary charge carrier in ZrNiSn,in this work we investigated the effects of Yb (4f 146s2) doping at the Zr (4d25s2) sites onthe TE properties of ZrNiSn.

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2. Experimental procedures

One general concern of HH alloy synthesis is the homogeneity of samples, and, as such,most synthesis procedures involve a long annealing time to help homogenize the sam-ple. To ensure sample homogeneity, avoid crucible-related cross-contamination and havea shorter synthesis time, we adopted a combined levitation melting and spark plasma sin-tering (SPS) procedure recently developed by our group [12]. Yb-doped Zr1−xYbxNiSn(x = 0, 0.01, 0.02, 0.04, 0.06 and 0.10) HH alloys were prepared by levitation melting ofstoichiometric amounts of Zr (99.99%), Ni (99.999%), Sn (99.999%) and Yb (99.999%)in an argon atmosphere for 2 min, and the melt was quenched in a water-cooled coppercrucible. The ingot was melted twice to ensure homogeneity. The ingots were pulverizedthrough a ball milling (BM) process at 200 rpm for 4 h. The powders were then pelletizedto ∼99% theoretical density using a SPS procedure (SPS-1050, Sumitomo Coal MiningCo.) at 1173 K for 10 min under 65 MPa.

The pellets were directly used to measure the thermal conductivity, and then cut intorectangular bars for electrical conductivity and Seebeck coefficient measurements. Thesemeasurements were carried out from 300 K to 900 K, if not otherwise noted. The phasestructure of the samples after BM and after SPS was checked by X-ray diffraction (XRD)on a Rigaku D/MAX-2550PC diffractometer using Cu Kα radiation (λ = 1.5406 Å).The electrical conductivity and the Seebeck coefficient were measured on a custom-designed, computer-assisted apparatus using a DC four-probe method and a differentialvoltage/temperature technique, respectively [29]. The thermal conductivity, κ , was calcu-lated by using κ = DρCp, where ρ is the sample density measured by an Archimedes’smethod. The thermal diffusivity, D, and specific heat, Cp, were measured by a laser flashmethod on a Netzsch LFA457 with a Pyroceram standard [30]. The uncertainties for D andCp measurements were ±3% and ±5%, respectively.

3. Results and discussion

The XRD patterns of the powders of levitation melted Zr1−xYbxNiSn alloys are shownin Figure 1a. All major peaks could be indexed to the ZrNiSn phase with MgAgAs-typestructure (JCPDS No.23-1281), while some weak peaks from a small amount of Sn weredetected. After the SPS process, all samples were single phased except the x = 0.1 sample,which showed a trace of impurity phase (Figure 1b). As shown in the inset of Figure 1b,the lattice constant increases with the increase of the nominal Yb doping ratio, and thevariation agrees fairly well with Vegard’s law. In view of the fact that the ionic radius ofYb (0.86 Å) is larger than that of Zr (0.79Å), this is direct evidence of Yb doping intothe lattice and indicates that the actual Yb doping ratio is proportional to the nominalone. These observations are also consistent with our previous work [27], showing that thecombined levitation melting and SPS procedure is effective in preparing high-quality HHalloy samples in a time-efficient manner.

The temperature dependences of electrical conductivity, σ , and Seebeck coefficient, α,of Zr1−xYbxNiSn (x = 0, 0.01, 0.02, 0.04, 0.06 and 0.10) alloys are shown in Figures 2aand b, respectively. The σ of all samples increases with increasing temperature, showingsemiconducting behavior. We note that Yb doping consistently decreases the magnitudeof σ over the entire temperature range studied. On the other hand, Yb doping impacts themagnitude and the temperature dependence of α in a non-monotonic fashion. The mag-nitude of α first increases with increasing Yb content up to x = 0.02 with a maximumα ∼ –210 µVK−1, above which it starts decreasing. In particular, the sign of α is changed

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10 20 30 40 50 60 70 80

(331

) (420

)

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)

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x = 0.06x = 0.04x = 0.02x = 0.01

Inte

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(a.

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2θ (deg.)

x = 0

Sn(a)

(200

) (220

)

(hkl) : ZrNiSn

10 20 30 40 50 60 70 80

0.00 0.02 0.04 0.066.1236.1246.1256.1266.1276.1286.1296.1306.1316.1326.1336.134

x = 0.10x = 0.06x = 0.04x = 0.02x = 0.01

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nsity

(a.

u.)

2θ (deg.)

x = 0

Impurity

(b)

Lat

tice

Con

stan

t (Å

)

Yb Content, x

Figure 1. XRD patterns of the Zr1−xYbxNiSn (x = 0, 0.01, 0.02, 0.04, 0.06 and 0.10) compoundsafter BM (a) and after SPS (b). The calculated lattice constants after SPS are shown in (b) as afunction of the nominal Yb doping ratio.

to positive (p-type) at room temperature for the x = 0.10 sample (see inset of Figure 2b),suggesting that heavy Yb doping establishes p-type conduction in the material althoughthe magnitude of α is not yet sufficient for a high p-type performance. The trend of thisn-type to p-type changeover is apparent for the x = 0.04, 0.06 and 0.10 samples near roomtemperature: the value of α changes from –61 µVK−1 to 14 µVK−1. In view of the elec-tronic configurations of Yb (4f 146s2), Y (4d15s2) and Zr (4d25s2), the sign change of α

suggests that Yb doping introduces acceptors into the material in the same way that Y dop-ing does [24,27]. In this context, the decrease of σ upon Yb doping can be thus ascribed tothe decreased carrier concentration due to the charge (electron–hole) compensation effect.The temperature dependence of α for all samples is typical of a degenerate semiconductor:the magnitude of α first increases with increasing temperature until reaching a maximum(inflection point), and then decreases due to the bipolar effect (Figure 2b). Leaving the mag-nitude of α aside, the maximum of α systematically shifts toward higher temperature withincreasing Yb content, indicating that the Yb doping alters the band structure of ZrNiSn.

Figure 3a shows the temperature dependence of the thermal conductivity, κ , of theZr1−xYbxNiSn alloys. In general, the κ is the sum of the carrier thermal conductiv-ity, κe, and the lattice thermal conductivity, κL, i.e. κ = κL + κe. The κe can beestimated by the Wiedemann–Franz relationship, κe = LσT with the Lorentz number

300 400 500 600 700 800 9000

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10x = 0x = 0.01x = 0.02x = 0.04x = 0.06x = 0.10

σ (1

04 Ω

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α (μ

V K

−1 )

T (K)

(b)

α (µ

V K

−1)

Yb Content, x

Figure 2. Temperature dependences of electrical conductivity σ (a) and Seebeck coefficient α(b) for Zr1−xYbxNiSn (x = 0, 0.01, 0.02, 0.04, 0.06 and 0.10) compounds. The room temperatureSeebeck coefficient α as a function of the nominal Yb-doping ratio is shown in the inset of (b).

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300 400 500 600 700 800 9004.5

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1 )

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x=0x=0.01x=0.02x=0.04x=0.06x=0.10

(a)

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Cp (J

g−1 K

−1 )

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0

2

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κ L(W

m−

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m−

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x = 0x = 0.01x = 0.02x = 0.04x = 0.06x = 0.10

x = 0x = 0.01x = 0.02x = 0.04x = 0.06x = 0.10

(b)

Figure 3. Temperature dependences of thermal conductivity k (a) and the electrical thermal con-ductivity ke and lattice thermal conductivity kL (b) for Zr1−xYbxNiSn (x = 0, 0.01, 0.02, 0.04, 0.06and 0.10) compounds. Temperature dependence of specific heat Cp is presented in the inset of (a).

0.004.5

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0.02 0.04 0.06

0.00360

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420

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θ D (

K)

0.08 0.10Yb Content, x

κ an

d κ L

(W

m−

1 k−

1 )

κL

κ

θD

Figure 4. Variation of k and kL with the Yb doping ratio at room temperature. Yb dopingdependence of θD is presented in the inset.

L = 2.0 × 10−8 V2 K−2. The calculated κL and κe are shown in Figure 3b. We note that(i) the κL significantly decreases while the κe only slightly decreases with increasing Ybdoping ratio, and (ii) the κL predominates the κ in all compounds. It is also noted that Ybdoping effectively reduce κ and κL in the vicinity of room temperature (Figure 4) but theextent of reduction quickly gets marginal at elevated temperatures. To elucidate the originof this observation, we inspect the temperature dependence of specific heat, Cp (inset ofFigure 3a) from 300 K to 900 K. It is plausible to assume that Cp is predominated by thelattice specific heat above room temperature. Furthermore, fitting the observed Cp by the

Debye formula of lattice specific heat, Cp = 9R(T/θD)3∫ θD/T

0 exp(x)x4/(exp(x) − 1)2dx,x = hω/kBT , ω is the frequency of phonons, yields the Debye temperature, θD, where Cp

is the molar specific heat. As shown in the inset of Figure 4, θD systematically decreaseswith increasing Yb doping ratio. We note that the derived θD value is somewhat higherthan previously reported [31], but this will not affect the following discussion as we areonly concerned with the trend of θD on doping. In a simple kinetic model, the κL is propor-tional to the product of Cp, the velocity of sound, θ s, and the phonon mean free path, lph.The value of θD is a rough gauge of the velocity of sound υs. In this vein, the decrease ofκL can be ascribed to enhanced phonon scattering and thus a decreased lph by the dopantinduced mass fluctuation mechanism.

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300 400 500 600 700 800 9000.0

0.5

1.0

1.5

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2.5

Pow

er F

acto

r (1

0−3

W m

−1

K−

2 )

T (K)

x = 0x= 0.01 x = 0.02x = 0.04x = 0.06x = 0.10

(a)

300 400 500 600 700 800 900

0.0

0.1

0.2

0.3

0.4

Figu

re o

f m

erit

ZT

T (K)

x = 0x = 0.01x = 0.02x = 0.04x = 0.06x = 0.10

(b)

Figure 5. Temperature dependences of power factor (a) and figure of merit ZT (b) for Zr1−xYbxNiSn(x = 0, 0.01, 0.02, 0.04, 0.06 and 0.10) compounds.

Finally, the temperature dependences of the power factor and figure of merit, ZT , ofZr1−xYbxNiSn alloys are shown in Figures 5a and b, respectively. Due to the degradedpower factor upon Yb doping, there is not much improvement in the ZT values of Yb-doped samples as compared with the pristine ZrNiSn. Nonetheless, further improvementof ZT may be attained by selectively doping the Ni and Sn sites to increase the power factor,and/or via the ‘nanocompositing’ approach recently reported by Yan et al. [32].

4. Conclusions

In this paper, we have reported the thermoelectric properties of Zr1−xYbxNiSn (x = 0, 0.01,0.02, 0.04, 0.06, 0.10) compounds prepared by a time-efficient levitation melting and sparkplasma sintering procedure. Yb doping appeared to introduce holes into the system, and theroom temperature Seebeck coefficient changed into p-type at a nominal Yb doping contentat x = 0.10 although the magnitude of Seebeck coefficient was not high. We did not seesignificant improvement of ZT by Yb doping. Further efforts, e.g. doping on the Ni and Snsites, will be carried out to further improve the ZT .

AcknowledgementsThis work is supported by the Nature Science Foundation of China (Grant Nos. 51171171 and50971115), and the National Basic Research Program of China (Grant No.2007CB607502). J.H.would like to thank the support by the National Science Foundation of the United States (GrantNo. 1008073).

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